Surface segregation in MCrAlY alloys

Surface Segregation in MCrAIY Alloys

C.L. BRIANT and K. L. LUTHRA

This paper presents a study of surface segregation in NiCrA1Y and FeCrA1Y type alloys. The segregation was measured using Auger electron spectroscopy. Samples were heated in the spectrometer to temperatures between 800 and 1100 ~ and segregation was measured at temperature as a function of time. The results show that when yttrium and sulfur were present in the alloy, they segregated to the surface where their concentrations were greatly enriched over their bulk concentrations. The presence of yttrium in the alloy did not eliminate sulfur segregation. The surface concentration of aluminum appeared, in most cases, to be greater than its bulk concentration. However, because of uncertainty in sensitivity factors for Auger emission, this point could not be conclusively proven. Cobalt and chromium, in contrast, were depleted from the surface. Sulfur segregation was also examined in Ni-S, Cr-S, A1-S, and Y-S binary alloys. Sulfur was found to segregate extensively in the nickel and chromium alloys, but to a much less extent in aluminum. No sulfur segregation was observed in the yttrium alloy.

I. INTRODUCTION

OXIDATION of iron and nickel base alloys has been stud- ied for many years because it often limits applications of these alloys at high temperatures. As a result of this research several statements can be made. One is that chromium and aluminum additions to the alloy help provide protection against oxidation of the base metal through the formation of Cr203 and AI~O 3 on the surface. Of these two oxides alumina is preferred because Cr20 3 is subject to volatilization above approximately 1000 ~ Another fact is that yttrium and hafnium additions to these allo~s reduce spallation of the oxide when the metal is cooled.I I

Many mechanisms have been proposed to explain why the additions of yttrium and hafnium improve adherence of the oxide, and they are summarized in Reference 2. None of these models is consistent with all of the experimental data. Recently, several researchers 13-6j have proposed that these two elements have a beneficial effect because they re- tard sulfur segregation to the metal surface through the formation of yttrium and hafnium sulfides in the matrix. They point out that, since sulfur is a strong embrittler of grain boundaries in iron and nickel base alloys, IT-l~ it might also weaken the metal-oxide bond if it is present on the metal surface. Therefore, in this model spallation is primarily attributed to the presence of sulfur at the oxide-metal interface. More generally, this proposal would suggest that spallation resistance depends on the composition of the alloy surface rather than the composition of the bulk alloy.

As a result of this proposal it has become essential to study surface segregation in these alloys. In this paper we report a study of segregation in NiCrAI and FeCrAI type alloys. We have measured this segregation by Auger electron spectroscopy. The results will show that sulfur and yttrium are clearly enriched on the surface and that cobalt and chromium are depleted. Aluminum appeared to be slightly segregated in most cases, but because Auger sensitivity factors are rather uncertain, one must be careful in interpreting this result.

C.L. BRIANT and K, L. LUTHRA are Staff Scientists with General Electric Company, Research and Development Center, P.O. Box 8, Sche- nectady. NY 12301.

Manuscript submitted May 27, 1987.

II. EXPERIMENTAL

The MCrAI(Y) alloys used in this study were fabricated from the metals of various purities listed in Table I.* The

*Unless otherwise indicated, the alloys are of commercial purity. The alloy compositions are expressed in weight percent, and the surface compositions are expressed in atomic percent.

alloys were prepared by casting in graphite or copper molds and were subsequently machined to produce samples of the desired size. In addition to the MCrA1Y type alloys we also prepared samples of the metals Ni, Cr, A1, and Y containing 100 ppm S.

We used Auger electron spectroscopy to monitor the composition of the alloy surfaces at temperatures similar to those used in the oxidation experiments. For the Auger experiments, the sample was ground on silicon carbide paper to a thickness of approximately 0.025" (0.625 m) and then ultrasonically cleaned. The sample was placed in a holder such as the one shown in Figure I, and the entire assembly was placed in the Auger spectrometer. The sample was heated in situ using tungsten wires, 0.007" (0.125 mm) in diameter, that were spot welded to it at opposite points. The temperature was measured using a chromel-alumel thermocouple spot welded to the back of the sample.

To start an experiment, the system was pumped to a background pressure of approximately 2 • 10-1~ torr. It was then back-filled with argon to a pressure of 5 • 10 -5 torr, and the sample surface was cleaned by argon ion sputtering. This cleaning was performed to avoid any competitive interactions between the segregating species and contami- nants deposited from the atmosphere or during handling. The initial cleaning usually required three to five hours. The spectrometer was re-evacuated and the sample was then heated to the desired temperature. Data were taken at temperature as a function of time. Repeated runs were made by cooling the sample to room temperature, sputter cleaning the surface with argon ions, and reheating.

We found that the Auger results depended on the exact way we ran the experiment. In particular we noted that the initial run on the sample after it was first placed in the spectrometer could give very misleading results. This point has been noted in previous work t"'lzl and is usually attributed to rapid diffusion of the segregating species along dislocations

METALLURGICAL TRANSACTIONS A VOLUME 19A, AUGUST 1988--2099

Table I. Elements Used in Preparation of MCrAIY Alloys Element Source Nominal Purity Used in Samples Designated

270 nickel Huntington Alloys 99.98 C.P., H. E Spectroscopic nickel Johnson Matthey 99.999 V.H.P. Electrolytic cobalt Afrimet Indusa 99.91 C.P., H. P. Cobalt rod Johnson Matthey 99.998 V.H.P. A-101 iron SCM Metal Products 99.99 C.P. Spectroscopic iron Johnson Matthey 99.9985 H.P., V. H. P. EL chrome VG (Cr) Elken Metals Co. 99.63 C.P. Iochrome (Cr) Materials Res. Corp. 99.999 H.P. Chromium pellets Johnson Matthey 99.999 V.H.P. Aluminum Alcoa 99.99 C.P., H. P. Aluminum ingot Johnson Matthey 99.9999 V.H.P. Yttrium ingot Research Chemicals 99.* C.P., H. P. Yttrium lumps Johnson Matthey 99.99 V.H.P.

C. E--Commercial Purity; H. E - -High Purity; V. H. P.--Very High Purity. *Primarily tantalum impurity.

TO POWER SOURCE

TUNGSTEN W l R E ~

~,.,,"'~ CERA MIC BLOCK

HIGH CONDUCTIVITY

PLATES

Fig. 1 - A schematic drawing of the heating stage used for the Auger electron spectroscopy experiments.

and grain boundaries. Once these sources are depleted, they will no longer contribute to the segregation. We also found that the temperature at which sputtering was performed was important. If sputtering was carried out at the temperature of interest and the run did not involve a heat-up period, the results were quantitatively, though not qualitatively, different from those presented here. The exact cause for this is not known, although it has been noted by other researchers, t13] We chose to use our particular program because we felt it was more like the oxidation experiments.

One problem with Auger electron spectroscopy is the dif- ficulty in quantifying the results. In simple alloys where only a few elements segregate, one can successfully use peak height ratios as a relative measure of segregation. However, in multi-component systems such as the ones used in this study peak height ratios are often not sufficient, and one must attempt to convert peak heights to atomic composition. In this study, this conversion was made following the procedure outlined in the Handbook for Auger Electron Spectroscopy. The Auger peaks and sensitivity factors (S.E) that we used were the following: A1 62 eV (S.E = 0.22), Y 127 eV (S.E = 0.134), S 150 eV (S.E = 0 . 7 6 5 ) , Cr 489 eV ( S . F . = 0 . 3 4 4 ) , Co 775 eV (S.F. = 0.231), Fe 703 eV (S.F. = 0.226), Ni 848 eV (S.E = 0.27), and AI 139 eV (S.E = 0.067). Admittedly these sensitivity factors are the simplest ones that can be used and do not include various matrix correction terms.

However, we have chosen to use these simple sensitivity factors for the following reasons. (1) Calculation of matrix- effect corrections, according to the procedure of Seah, t141 suggests that these effects are not great. For example, for sulfur segregation in nickel, chromium, and aluminum the matrix corrections would be 1.15, 1.22, and 1.76, respectively, times our reported value. For yttrium segregation in nickel the matrix correction would be 0.71 times our values. (2) There is considerable debate about the correct values to be used for the various factors that go into matrix corrections, t~4j Therefore, one cannot be certain that by using them one has a more correct value for concentrations. (3) With the exception of aluminum segregation, all of the changes reported in this paper are very large. Since our conclusions are based on these large changes rather than on small differences in concentrations, we feel that they will not be affected by the particular correction scheme that we use. (4) By using the simple sensitivity factors listed above, one can easily work back to the original peak heights and treat the data with any scheme desired. (5) Since the NiCrA1 and NiCrAY alloys used for the major part of this investigation all have similar bulk compositions, the matrix corrections should be virtually identical for each sample, and therefore relative changes measured in this study should accurately reflect changes in the surface composition.

Transmission electron microscopy was used to examine the microstructure of the near surface region of samples that had been heated in the Auger spectrometer. The samples were electropolished on one side only in 20 pct perchloric acid in methanol. By polishing from only one side we were able to preserve the surface on which segregation had occurred. For comparison, a sample was also electropolished from both sides to reveal the bulk microstructure. The samples were analyzed in a Hitachi H-600-1 analytical electron microscope that operates at 10 kV.

III. RESULTS

A. Auger Electron Spectroscopy

Isothermal Aging at 1100 ~ We will begin by present- ing a complete description of the data obtained on a sample of commercial purity Ni-23Co-20Cr-12.5AI-0.3Y. These

2100--VOLUME 19A, AUGUST 1988 METALLURGICAL TRANSACTIONS A

results will p rovide a basel ine for compar i son with results for o ther alloys. F igure 2 shows the a l u m i n u m concen t ra - t ion on the surface plotted as a funct ion o f ag ing t ime for this commerc i a l purity alloy. Note that data are inc luded for six different runs for one A u g e r sample . The surface alu- m inum concentra t ions are very near the bulk concent ra t ion . Also , note that the equi l ibr ium concent ra t ion is es tabl ished in the first three to f ive minutes o f aging. The resul ts for a luminum in this al loy appear to be reproduc ib le f rom run to run on a g iven sample , but there is some var iabi l i ty f rom sample to sample as shown by the entr ies in Tab le II denoted by an asterisk. In this second sample the surface alu- m inum concentra t ion appears to be above the bulk level .

F igure 3 shows a s imilar plot for yt t r ium. Note that the values obtained during the first run are s igni f icant ly lower than those obtained in addit ional runs. Also , in the first run approximate ly one hour was requi red to reach equ i l ib r ium, whereas in subsequent runs there was no apparent change with t ime after the first few minutes o f ag ing . The re was some scatter in the data f rom run to run wi th the apparent e q u i l i b r i u m c o n c e n t r a t i o n v a r y i n g b e t w e e n abou t 5 and 10 at. pct. However , in all o f these runs the concent ra t ion o f yt t r ium on the surface was be low detec t ion after sputtering. Therefore , the observed yt t r ium levels were a resul t o f segregat ion at 1100 ~

0.5

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--" 0.1

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�9 RUN 6 o �9 .~,

0 I I I l [ t t t 1 [ 1 I I I ] I I I I

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Fig. 2 - - T h e aluminum concentration on the surface plotted as a function of heating time. The heating temperature was ll00 ~ and the alloy was the commercial purity Ni-23Co-20Cr-12.5AI-0.3Y.

Figure 4 shows a s imi la r p lo t for the change in sul fur concentra t ion with aging t ime. In each run, sulfur requires l o n g e r than y t t r i um and a l u m i n u m to app roach its equ i - l i b r ium va lue , and in the first run the t ime is no t i ceab ly longer than in the others. However , in each case the surface concentra t ion approaches a value o f about 10 at. pct.

Table II. Surface Compositions Found on Samples Equilibrated at I100 ~

Alloy Composition Surface Composition (At. Pct)

Wt Pct (At. Pct) Ni Fe Co Cr A1 Y P S

Ni23Co20Cr 12.5A10.3Y 40.4 - - 6.9 2.4 26.9 10.9 0 11.3 (Ni 19Co 19Cr23AI0.15Y) 54.7 - - 5.5 0 21.4 8.2 0 9.9

61.1 - - 4.5 1.5 16 7.7 0 8.9 57.9 - - 6.0 1.1 17.7 7.4 0 9.6 60.9 - - 5.9 0.8 16.8 5.9 0 9.4 61.7 - - 6.4 2.1 16.3 3.3 0 7.4

*42.2 - - 5.5 0.57 29.4 8.8 0.6 11.6 *43.5 - - 4.5 0.76 31.1 7.8 5.5 11.6

Ni20Crl2.5AI0.3Y 45 1 34.4 6.4 - - 13.2 0.03S (NiI9Cr23AI0.15Y0.045S) 46.9 - - - - - - 30.2 3.5 - - 19.4 high purity 50.7 - - - - 0.5 24.3 5.1 - - 19.4

41.8 - - - - 0.6 29.1 6.3 - - 22.2

Ni20Crl20.5AIO.3Y 53.6 - - - - - - 21.7 12.5 - - 12.2 (Ni19Cr23A10.15Y) 46.8 - - - - 1.8 28.7 17.2 - - 5.5 high purity

Ni20Crl2.5AI** 49.2 - - - - 0.96 40.2 - - - - 9.6 (Ni19Cr23AI) 50.2 - - - - 4.6 35.4 - - - - 9.8 high purity

Ni20Crl2.5AI 53.9 - - - - < 1 45.1 - - (Ni19Cr23A1) 54.5 - - - - 1.01 43.9 - - very high purity 53.6 - - - - 0.78 44.6 - -

59.0 - - - - 0.65 40.2 - - 57.7 - - - - 0 42.2 - -

Fe20Crl2.5A10.3Y - - 46.1 - - 8.7 20.6 12.4 (Fe 19Cr22A10.15Y)

Fe20Crl2.5A1 - - 57.8 - - 8.4 31.0 - - (Fe19Cr22A1)

Fe25Cr4AI0.3Y) - - 53.9 - - 15.9 24.6 0.6 (Fe25Cr8A10.15Y)

*Data obtained from a second Auger sample. **In our earlier paper, ltr; this sample was erroneously indicated as C. P. (commercial purity).

1.4

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O o

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TIME, hrs.

Fig. 3 - - T h e yttrium concentration on the surface plotted as a function of heating time. The heating temperature was 1100 ~ and the alloy was the commercial purity Ni-23Co-20Cr-12.5Al-0.3Y alloy.

0.12

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TIME, hrs. Fig. 4 - - T h e sulfur concentration on the surface plotted as a function of heating time. The heating temperature was It00 ~ and the alloy was the commercial purity Ni-23Co-20Cr-12.5AI-0.3Y alloy.

The final two elements that we wish to consider are cobalt and chromium. During any given run on either of the two samples used the concentration of cobalt on the surface would decrease from its bulk value to a value between 5 and 7 at. pct. However, sputtering would bring it back to a value near its bulk level. Chromium was also depleted but, in contrast to cobalt, sputtering for up to three hours would never return its surface concentration to the bulk value, and with repeated runs the concentration would drop to increas- ingly lower values. This point is shown graphically in Figure 5, where each break in the curve corresponds to a point where the sample was cooled and cleaned by sputtering. Based on calculations that included the partial pressure of chromium and its evaporation rate in vacuum, we would suggest that this depletion of chromium could be caused by evaporation into the vacuum chamber.

Table II provides a summary of the equilibrium values (i.e., the values at the end of a given run) for the above NiCoCrAIY alloy and the other alloys used in this study. In many cases, results are shown for more than one run to indicate the variability in experimental results; more runs than those listed were made on each sample. For all of the

12

10

8

, r ~ ~ ~1

.01

~ O

0.1 I.'0 - "10. TIME {hrs )

100.0

Fig. 5 - - The change in the surface concentration of chromium through a series of heating cycles at 1100 ~ on a sample of the commercial purity Ni-23Co-20Cr-12.5AI-0.3Y. Each break in the curve represents a point where the sample was cooled to room temperature and sputtered clean with argon ions.

alloys the following general comments can be made. Yt- trium and sulfur, when they are present in the alloy, are enriched on the surface over their bulk concentrations; chromium and cobalt are depleted. The aluminum concentration appears to range from near the bulk level to about twice the bulk level. Because of the difficulties described above in quantifying Auger data, it is difficult to be abso- lutely certain that aluminum is significantly enriched on the surface. There are, however some significant quantitative differences among the resu]ts listed in Table I1 for the different alloys, and we now discuss them.

Consider first the sulfur-doped NiCrAIY alloy which is the second alloy listed in Table II. This alloy, which is designated by its nominal composition of 0.03 wt pct sulfur, actually contained 0.023 wt pct sulfur. The surface aluminum concentration appears to be slightly higher than that observed in the NiCoCrAIY alloy discussed above, and the yttrium concentration appears to be slightly lower. How- ever, the most significant difference is in the amount of sulfur segregation. This difference is emphasized in Figure 6 where we have plotted the results for two runs on this sample. If one compares the results for this material with those for the NiCoCrA1Y alloy shown in Figure 4, it is clear that sulfur segregation is much greater in this sulfur-doped alloy.

The results for the third alloy listed in Table I! are also plotted in Figure 6. This alloy was made from the same high purity starting materials used to make the sulfur doped alloy. The results in Figure 6 show that the sulfur segregation is much less than that observed in the sulfur doped alloy but comparable to that observed in the commercial purity NiCoCrAIY alloy. The aluminum concentration is also similar to that observed for the above two alloys. The amount of yttrium on the surface of this alloy was somewhat higher than that found for the NiCoCrAIY alloy and significantly higher than that found on the surface of the sulfur doped alloy.

The fourth and fifth entries in Table II are for alloys that contained Ni, Cr, and A1 but not Y. The alloy in the fourth


z 0

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Fig. 6 - - T h e sulfur concentration on the surface plotted as a function of heating time for the sulfur doped NiCrAIY alloy and a NiCrAIY alloy made from high purity starting materials. The temperature of the samples was 1100 ~

entry was made from the same purity elements used to make NiCrA1Y alloys in the second and third entries, and the one in the fifth entry was made with the highest purity materials available to us. (In our earlier paper I161 the alloy in the fourth entry was erroneously indicated as commercial purity. The only difference between the alloys of high purity and commercial purity was in the purity of Cr used, as given in Table I.) Several important points should be noted. Chromium was again significantly decreased below its bulk value, but the enrichment of aluminum appeared to be greater than in any of the yttrium containing alloys. In the high purity material the sulfur concentration was similar to that observed in the other commercial purity and high purity nickel-based alloys. However, for the alloy made from the highest purity starting materials sulfur segregation was virtually eliminated. This extreme difference is shown in Figure 7. In order to check this point more completely, several different samples were cut from the ingot and examined. In each case we observed identical behavior. During the initial heating cycle a small amount of sulfur segregation was observed. However, after that sulfur was sputtered away, only very small amounts were observed during any additional heating cycles. Therefore, in order to eliminate sulfur segregation one must employ the extremely high purity starting materials used for this alloy.

To evaluate quantitatively the segregation factor (i.e., the ratio of sulfur concentration at surface to that in the bulk), we measured the bulk sulfur levels of various nickel- base alloys. Two techniques were employed: (i) a LECO in- strument, in which a known quantity of the sample is combusted, and the gases are analyzed by I R absorption, and (ii) Glow Discharge Mass Spectrometry. The latter technique was used on our samples by Charles Evans &

120

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0001 0.01 0.1 1.0 l0 0 TIME (hrsl

Fig. 7 - - T h e surface concentration of sulfur plotted as a function of heating time for samples of the high purity NiCrAI alloy and the NiCrAI alloy made from the highest purity starting materials available for use. Also included are data from the NiCrAIY alloy made from high purity starting materials. The temperature of the runs was 1100 ~

Associates in Redwood City, California. For commercial purity NiCrAIY samples, the sulfur level was observed to be of the order of 15 ppm by weight. For high and very high purity NiCrAI and NiCrAIY samples, the sulfur levels were not reproducible, even for the different sections of the same casting. They generally ranged from a low of about 1 ppm (by weight) to a maximum of about 8 ppm (by weight), except in one sample of a high purity NiCrA1 sample which showed 14 ppm (by weight). (Another section of the same casting indicated 1 ppm sulfur.) If one assumes a sulfur level of ~ 8 ppm (atomic fraction ~ 1.2 • 1-5), the segregation factors are of the order of 4500 to 10,000 for high purity NiCrA1Y and about 8000 for high purity NiCrA1.

Table II also shows results for iron-based alloys. In most respects these results are quite similar to those for the nickel-base alloys. Enrichment of both yttrium and aluminum was observed at the surface. Sulfur segregation was also observed, but we consistently found that the amount on the surface was less than that observed on the nickel-base alloys. Also, the amount of chromium depleted from the surfaces of these alloys did not appear to be as great as in the nickel-base materials.

The other impurity element that was observed to segregate on some samples was phosphorus. Those samples where it was observed are noted in Table II. The fact that it was not observed in other samples simply suggests that the bulk concentration was too low for segregation to be observed. When segregation did occur at 1100 ~ it was very rapid. With time, however, the phosphorus signal would decrease as sulfur began to segregate to the surface. There appears to be a competition for sites on the surface between these two elements, as has been noted in surface segregation studies on other alloys. Illl

The final point to consider is the length of these runs at 1100 ~ Most runs lasted between one and two hours, because it was clear that the equilibrium composition was established very rapidly. However, to be certain that this was sufficient time we held one sample of a NiCrA1Y alloy at 1100 ~ for seven hours. We found that equilibrium was rapidly established and did not change throughout this run.

Segregation to Quenched Surfaces. We measured the change in segregation that occurred during cooling to room


temperature. At the end of each run at 1100 ~ the sample was cooled as rapidly as possible by cutting off the power supply that heated the tungsten wires. In general the sample cooled to below 350 ~ in about 30 seconds and to below 100 ~ in about 3 minutes. Surface compositions obtained from these quenched samples are given in Table III. The main effect of this cooling treatment was to produce a slight increase in the yttrium concentration and a significant increase in the sulfur concentration. Note that in the sulfur doped alloy the concentration of sulfur on the surface is ex- ceptionally high, but that in the highest purity NiCrA1 sam- pie sulfur is still absent.

Effect of Temperature on Aging. A few samples were heated to temperatures between 800 and 1000 ~ to study the effect of temperature on segregation. Figures 8 and 9 show the temperature dependence of segregation for aluminum and yttrium, respectively. Although in the case of aluminum there may be a slight decrease in the surface concentration with increasing temperature, neither element

z 0 p - t j

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0.45

0.40

0.35

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0.20 I I I 800 900 1000 tt O0

TEMPERATURE (%)

Fig. B--The temperature dependence of the surface aluminum concentration.

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TEMPERATURE (*C) Fig. 9--The temperature dependence of the surface yttrium concentration.

shows a strong temperature dependence in its tendency to segregate. All values in these two figures represent the surface concentration after one hour of aging, which in all cases is near the equilibrium concentration.

The results for sulfur segregation are somewhat different. Figure 10 shows the change in sulfur on the surface plotted as a function of time for four different aging temperatures. All data were taken on the high purity NiCrAI sample. As the temperature increases, less time is required to achieve equilibrium. Although the differences observed among the data taken at 900, 1000, and l l00 ~ may be purely a diffusion-controlled effect, the causes for the significant decrease in segregation kinetics at 800 ~ are not clear. Identical results were obtained for the commercial purity NiCoCrAIY sample.

Effect of Oxygen on Segregation. The results described so far were all obtained in a high vacuum environment. Since one of our goals is to relate this work to oxidation, we wanted to determine if these results would also be observed in an oxygen containing environment. Two kinds of experiments were conducted to answer this question.

Table I lL Surface Composit ions Found on Samples Equilibrated at 1100 ~ and Then Quenched to R o o m Temperature

Elemental Composition (At. Pct)

Alloy Ni Fe Co Cr A! Y P S

Ni23Co20Cr12.5A1 42.0 4.3 1.3 29.8 2.0 0.3Y 43.7 4.5 1.3 31.7 5.6

Ni20Crl2.5AI0.3Y 36.4 - - - - 2.1 29.2 6.0 0.03S 36.6 - - - - 0.7 25.6 6.9 high purity

Ni20Cr 12.5A10.3Y 55.5 - - - - 6.1 20.6 3.2 high purity 42.1 - - - - - - 21.6 14.3

Ni20Crl2.5AI 56.7 - - - - - - 43.3 - -

Ni20CrI2.5A1 43 7.3 31.2 high purity

7.3 12.9 0 13.0

- - 26.3 - - 30.2

B

m

B

14.2 22.0

18.5


0t4 . . . . . . . . I . . . . . . . . I . . . . . . . . I i , i i . . . .

o = 0.12- / " A

o.,0- " / , / I o 800'C

008- / /~ r o 9000C o = / / / ,, ooooc o.oo- / / / [],,00oc

-o.o,_ / / 1 �9 . t00

TIME (hrs) Fig. 1 0 - - T h e surface sulfur concentration plotted as a function of heating time for four different temperatures. These data were obtained on the high purity NiCrA1 sample.

A sample of high purity Ni-20Cr-12.hAI was sputter cleaned under high vacuum conditions in the normal fash- ion. Air was then admitted to the system until the pressure had reached one atmosphere. The spectrometer was then evacuated to a pressure of 3 x 10 -7 torr. The sample was heated to approximately 1050 ~ for ten minutes and quenched. Figure l l(a) shows a spectrum obtained from this sample and Figure ll(b) shows a spectrum from a sample with a similar heating cycle at 5 x 10-t~ torr. The oxy-

L I J

z

I J J

Z

I I I I I I I I I i | |

S At 159E eV

Ni 200 400 600 800 1 O0 0

ELECTRON ENERGY (eV) (a)

I I I I I I I I I I |1

At 1596eV

I I I { I I I I 1 Ill 0 200 400 600 800 t000

ELECTRON ENERGY (eV) (b)

Fig. 1 1 - Auger spectra obtained from the high purity NiCrAl saml;les. (a) The sample was heated in the spectrometer with a background pressure of 3 x 10 -7 toP/'. (b) The sample was heated in the spectrometer with a background pressure of 5 x 10 -1~ torr.

gen peak is clearly much greater in the sample heated at 3 x l 0 -7 tort, but all other peaks are similar. Therefore, we can conclude that oxygen at this pressure does not affect qualitatively the segregation results.

The final type of experiment performed to examine this issue was to oxidize a sample in air for very short times. The shorter times would hopefully keep the oxide film much thinner so that after short sputtering times one could obtain Auger data from the material near the oxide-metal interface. When such experiments were performed, we were able to detect a significant sulfur signal beneath the oxide film. This result also suggests that the sulfur segregation observed on the surfaces of these alloys in air is representative of that which occurs in Auger experiments performed in high vacuum.

Segregation in Binary Alloys. In addition to our study of MCrA1Y alloys, we examined segregation in Ni-S, Cr-S, A1-S, and Y-S binary alloys. The goal of this work was to determine if results obtained in these simple alloys would be similar to those obtained for the compositionally complex MCrA1Y materials. If they are similar, then we can assume that the thermodynamic equilibria that control segregation in the complex alloys are the same as those that control segregation in the simple alloys.

Table IV lists the results of these experiments. Signifi- cant sulfur segregation occurred in both the chromium and nickel alloys, al though the amount of segregation in chromium appears to be slightly less than in the nickel. In contrast when the aluminum-sulfur alloy was heated to 600 ~ which is a greater fraction of its melting point than 1000 ~ is of the melting point of nickel or chromium, very little sulfur segregation occurred. However, it should be noted that the surface of the aluminum sample always yielded very strong carbon and oxygen signals. These per- sisted even though the sample was repeatedly heated and sputtered. The yttrium-sulfur alloy showed no sulfur segregation at all.

B. Microstructural Characterization

One interesting point noted in the Auger results was the fact that chromium was significantly depleted from the sur-

Table IV. Sulfur Segregation in Binary Alloys

Temperature Sulfur Metal of Run (~ (At. Pct)

Ni 1000 16.3 1000 15.3 1000 15 1000 16.7

Cr 1000 13.5 1000 13 1000 10.1 1000 11.1

A1 600 1.5 600 1.7 600 0.95

Y ~ 1000" 0 1000" 0

*The approximate sign is used with these temperatures because we found it very difficult to obtain a good spot weld of the thermocouple onto the yttrium. Consequently, we had to use the current passing through the heating wires as a measure of temperature.


face but that aluminum was enriched. The depth of this depletion was greater than the depth of material removed in three hours of sputtering. To check this point in more detail we prepared thin foils for analytical electron microscopy, as described in the Experimental Section. This work was performed on both the highest purity NiCrA1 and commercial purity NiCoCrA1Y alloys. In both of these materials the surface region was found to be depleted of chromium, as shown in Table V. In the NiCrAI sample this surface material was found to have a lath-like structure, as shown in Figure 12(a). In the NiCoCrA1Y sample laths were not observed, as shown in Figure 12(b). Electron diffraction showed that in the NiCoCrA1Y alloy this surface phase was /3-martensite (NiAI). The diffraction patterns obtained from the NiCrA1 phase could not be unequivocably indexed, but were consistent with this phase also being fl-martensite.

For comparison we examined the bulk (i.e., nonsurface) microstructure of one of the Auger samples of the highest

purity NiCrA1 alloy. The structure was significantly different. It consisted of dendrites, denoted by D in Figure 13, and interdendritic regions, denoted by I in the figure. These latter regions thinned preferentially. Electron diffraction showed that the interdendritic regions contained two phases which were identified as a-Cr and 3/(Ni3(Cr, A1)). Typical compositions of these regions are given in Table V. The dendritic regions have compositions typical of fl-martensite but electron diffraction could not conclusively identify this phase. The dendrites also contained precipitates of a-Cr.

In addition to these studies we attempted to use analytical electron microscopy to examine the metal-oxide interface for surface segregation. However, we were unsuccessful in preparing thin foils that preserved this interface.

IV. DISCUSSION

We wish to begin this section by considering some general ideas regarding the surface segregation results. Sulfur has been reported to segregate to the surfaces of many alloys when they are heated. The results in this study show that sulfur segregation occurred quite readily in both MCrA1 and MCrA1Y alloys with the exception of the extremely

(a) (b) Fig. 12--Transmission electron micrographs of thin foils prepared from the material very near the surface from which the Auger data were obtained. The samples were (a) the highest purity NiCrAI sample and (b) the commercial purity NiCoCrA1Y sample.

Table V. Composition of Near Surface and Bulk Phases Obtained by Analytical Electron Microscopy

Composition (At. Pct)

Location Alloy Phase Ni Cr A1 Co

Surface NiCoCrA1Y fl-martensite 42 2.2 42.7 42.7 2.3 41.6

Surface NiCrA1 /3-martensite? 58.4 3.7 38 high purity 60.0 3.7 36.3

Bulk NiCrA1 ct-Cr 23.6 65.9 10.5 highest purity 3~' 62.6 23.5 13.9

/3-martensite? 59.8 6.9 34.3

13.1 13.4


Fig. 1 3 - A transmission electron micrograph of a thin foil prepared from the material well beneath the surface of an Auger sample of the highest purity NiCrA1 alloy.

high purity NiCrA1 sample. Smialek and Browning t6j and Smeggil et al.tSJ have also shown that sulfur segregates in both NiCrA1 and NiCrAIY alloys. The question that we must consider is whether in MCrA1 and MCrA1Y alloys this segregation is somewhat anomalous. Nickel, chromium, aluminum, and yttrium all form stable sulfides, with their stability increasing in the order in which the elements are listed. Furthermore, the results on the binary alloys show that for nickel, chromium, aluminum, and yttrium the amount of segregation appears to decrease as the stability of the sulfide increases. In particular, we found no sulfur segregation for yttrium. Therefore, it might seem surprising that we observe any sulfur segregation in the multicomponent alloys.

All of our segregation results are explained quantitatively by a thermodynamic model presented in another paper, t~51 Here we will present a summary of the qualitative argu- ments. Segregation equations that we have developed indicate that there are two important factors to be considered in segregation in these alloys. They are the following:

1. The first is the surface tension of the segregating element in comparison to the primary element. If the surface tension of the segregating element is less than that of the primary metal, it will tend to increase the segregation factor (the ratio of the surface to the bulk concentration) above unity. 2. The second is the activity coefficient of the segregating element. A negative deviation from ideality, representative of an activity coefficient of less than unity, would tend to reduce segregation.

Let us first consider the segregation of Cr, A1, and Y in NiCrA1Y type alloys. The surface tensions of chromium and nickel are comparable. The observed depletion in chromium could be caused by the fact that it has a small negative deviation from idealitys(i.e., by an activity coeffi-

I I cient value of less than unity). However, it could also be partly caused by evaporation, as mentioned above. The activity coefficients of aluminum and yttrium in NiCrAIY al-

loys are even less than that of chromium. However, their segregation is favored because their surface tensions are substantially lower than that of nickel.

Let us now consider the segregation of sulfur in the binary alloys. Its surface tension is much lower than that of any of the other elements--nickel, chromium, aluminum, and yttrium. Due to strong bonding of sulfur with these elements, its activity coefficient is also much lower than unity. The first factor is dominant for nickel, chromium, and aluminum. Consequently, sulfur segregation is observed in nickel-sulfur, chromium-sulfur , and aluminum-sulfur alloys. However, the sulfur bonding with yttrium is very strong and this effect dominates. Consequently, the second factor is more important, and segregation factors below unity are observed for yttrium-sulfur alloys. The segregation of sulfur on NiCrA1 and NiCrA1Y alloys is further complicated because of the presence of many elements. Higher segregation factors are observed on the surface of these alloys in comparison to aluminum-sulfur and yttrium- sulfur alloys for the following reasons:

1. Owing to the presence of chromium and nickel, the surface tensions of the NiCrAI and NiCrA1Y alloys are much higher than for aluminum and yttrium, respectively. 2. Owing to the formation of nickel aluminides (Ni3A1) and nickel yttrides (Ni2YlT), the aluminum and yttrium ac- tivities would be less than unity in these multicomponent alloys. Our calculations show that due to the low aluminum activity in NiCrA1, the aluminum sulfide is unstable. The most stable sulfides in NiCrA1 and NiCrA1Y alloys are CrS and Y2S3, respectively. The resulting activity and activity coefficients of sulfur in NiCrA1 and NiCrAIY alloys are much higher than in AI-S and Y-S binary systems, respectively.

Let us now consider the influence of segregation on oxidation, which, as stated in the Introduction, was our original reason for doing this work. Two important ways in which we believe segregation can influence oxidation pro- cesses are the following: (i) the scale/metal adhesion, and (ii) transient oxidation.

In one of our other papers, ~61 we have discussed the mechanism of scale/metal adhesion in depth. As summarized in this paper, we find little or no correlation between the sulfur segregation and the scale/metal adhesion. There- fore, we have concluded that sulfur segregation is not the primary cause of scale spallation. The two most important pieces of evidence that allow us to draw this conclusion are the facts reported in Reference 16 that extensive spallation occurred in the highest purity NiCrA1 alloy even though sulfur segregation was essentially eliminated and that additions of 0.03S to a NiCrA1Y alloy did not degrade its spallation resistance even though, as shown in Figure 6, it caused extensive segregation of S. We have also suggested that scale adhesion might be enhanced by yttrium segregation. E~rj Although conclusive evidence for this point is lack- ing, it would be consistent with the calculations reported by Johnson and Pepper. I~Tj

Let us now consider the effect of segregation on the early stages of oxidation. According to the current concepts of oxidation, when an alloy is heated in an oxidizing atmosphere, the surface is oxidized through a series of stages (referred to as transient oxidation) before achieving a final, steady state composition. The rate of transition from this


initial oxidation to a steady state scale would, among other factors, depend on the relative proportion of different oxides on the surface. The segregation results observed here suggest that the relative proportion of the oxides of different elements on the surface may be significantly different from the relative proportion of elements in the alloys. For example, on Ni-20Cr-12.5AI-0.3Y alloys, the initial scale would be expected to contain Cr203, NiO, and A1203. Since aluminum oxide is the most stable oxide, the final steady state scale would consist of aluminum oxide at the scale/metal interface. There would be a thin layer containing chromium oxide and nickel oxide, formed during the early stages, at the gas/scale interface. Our Auger studies indicate that the surface concentration of chromium is substantially lower than the bulk concentration. This would reduce the transition time from the transient oxidation to steady state oxidation. This transition time would also be influenced by the significant concentration of yttrium observed in our Auger studies.

V. CONCLUSIONS

The conclusions of this study are the following:

1. Sulfur segregates to surfaces of NiCrAIY and FeCrAIY alloys when they are heated at 1100 ~ The rate at which sulfur segregates decreases as the heating temperature decreases. When a sample is quenched to room temperature, sulfur segregation increases during the quench.

2. Yttrium is also enriched on the surfaces of these alloys during heating at 1100 ~ The sulfur and yttrium concentrations do not have a strong temperature dependence in the temperature range of 800 to 1100 ~

3. Chromium and cobalt are depleted from the surface as a result of heating. The depletion of chromium occurs slowly, but its bulk concentration cannot be regained even after sputtering for 3 hours.

4. Experiments on binary metal-sulfur alloys showed that sulfur readily segregates in nickel and chromium. Segre- gation is greatly decreased in aluminum and did not occur in yttrium.

5. When these results are combined with the oxidation re- suits presented in Reference 16, it is found that no correlation can be made between spallation resistance and sulfur segregation.

ACKNOWLEDGMENTS

The authors would like to thank Mr. Nathan Lewis and Dr. Ernest Hall for performing the transmission electron microscopy.

REFERENCES

1. L.B. Pfeil: U.K. Patent No. 459848, 1937. 2. D.P. Whittle and J. Stringer: Phil. Trans. Roy. Soc. London A,

1980, vol. 295, p. 309. 3. J.G. Smeggil, A.W. Funkenhnsch, and N. S. Bomstein: presented

at the spring 1984 meeting, Electrochemical Society, Cincinnati, Oil, May 1984.

4. K.L. Luthra and C. L. Briant: presented at the spring 1984 meeting, Electrochemical Society, Cincinnati, Oil, May 1984.

5. A.W. Funkenbusch, J.G. Smeggil, and N. S. Bornstein: Metall. Trans. A, 1985, vol. 16A, p. 1164.

6. J.L. Smialek and R. Browning: Proc. Symp. on High Temperature Materials II1, Z. A. Munir and D. E Cubicciotti, eds., Electrochem. Soc., Pennington, NJ, 1986, p. 258.

7. C.L. Briant: Acta Metall., 1985, vol. 33, p. 1241. 8. C. Pichard, J. Rieu, and C. Goux: Metall. Trans. A, 1976, vol. 7A,

p. 1811. 9. R.A. Mulford: Metall. Trans. A, 1983, vol. 14A, p. 865.

10. R.A. Mulford: in Embrittlement of Engineering Alloys, C.L. Briant and S.K. Banerji, eds., Academic Press, Treatise on Materials Sci- ence, 1983, vol. 25, p. 1.

11. C.L. Briant and R.A. Mulford: Metall. Trans. A, 1982, vol. 13A, p. 745.

12. J.J. Burton, B.J. Berkowitz, and R.D. Kane: Metall. Trans. A, 1979, vol. 10A, p. 677.

13. M. Oku, K. Hirokawa, H. Kimura, and S. Suzuki: Surface andlnter- face Analysis, 1986, vol. 9, p. 67.

14. M.P. Seah: in Practical Surface Analysis, D. Briggs and M. P. Seah, eds., John Wiley and Sons, Ltd., London, 1985, p. 181.

15. K.L. Luthra and C.L. Briant: Metall. Trans. A, 1988, vol. 19A, pp. 2091-98.

16. K.L. Luthra and C.L. Briant: Oxidation of Metals, 1986, vol. 26, p. 397.

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Surface segregation in MCrAlY alloys